CN114585766A - High-strength steel sheet and method for producing same - Google Patents

High-strength steel sheet and method for producing same Download PDF

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Publication number
CN114585766A
CN114585766A CN202080073477.6A CN202080073477A CN114585766A CN 114585766 A CN114585766 A CN 114585766A CN 202080073477 A CN202080073477 A CN 202080073477A CN 114585766 A CN114585766 A CN 114585766A
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less
mass
retained austenite
average
steel sheet
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CN202080073477.6A
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CN114585766B (en
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远藤一辉
川崎由康
田路勇树
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JFE Steel Corp
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JFE Steel Corp
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21C—MANUFACTURE OF METAL SHEETS, WIRE, RODS, TUBES OR PROFILES, OTHERWISE THAN BY ROLLING; AUXILIARY OPERATIONS USED IN CONNECTION WITH METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL
    • B21C47/00—Winding-up, coiling or winding-off metal wire, metal band or other flexible metal material characterised by features relevant to metal processing only
    • B21C47/02—Winding-up or coiling
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B32—LAYERED PRODUCTS
    • B32B—LAYERED PRODUCTS, i.e. PRODUCTS BUILT-UP OF STRATA OF FLAT OR NON-FLAT, e.g. CELLULAR OR HONEYCOMB, FORM
    • B32B15/00—Layered products comprising a layer of metal
    • B32B15/01—Layered products comprising a layer of metal all layers being exclusively metallic
    • B32B15/013—Layered products comprising a layer of metal all layers being exclusively metallic one layer being formed of an iron alloy or steel, another layer being formed of a metal other than iron or aluminium
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    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
    • C—CHEMISTRY; METALLURGY
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
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    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/78—Combined heat-treatments not provided for above
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/0236—Cold rolling
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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    • C22C—ALLOYS
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    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
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    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
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    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
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    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
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    • C23C2/06—Zinc or cadmium or alloys based thereon
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    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

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  • Heat Treatment Of Sheet Steel (AREA)

Abstract

A high-strength steel sheet having a diffusible hydrogen content of 0.3 ppm by mass or less in steel, having a predetermined composition, and having a steel structure comprising: the steel sheet has a structure in which ferrite is 30 to 80% in terms of area percentage, tempered martensite is 3.0 to 35% in terms of area percentage, retained austenite is 8 to 8%, the area percentage of retained austenite having an aspect ratio of 2.0 to 1 μm in terms of the minor axis is 0.3 or more, the area percentage of all the retained austenite is 1.5 or more, the average Mn amount (mass%) in the retained austenite is 3.0 or more, the average C amount (mass%) in the retained austenite is 3.0 or more, and the average C amount (mass%) in the retained austenite is 3.0 or more, the average C amount (mass%) in the retained austenite is small, the average Mn amount (mass%) in the retained austenite is small At 0.05.

Description

High-strength steel sheet and method for producing same
Technical Field
The present invention relates to a high-strength steel sheet having excellent formability and a method for producing the same, which is suitable as a member used in the industrial fields of automobiles, electric appliances, and the like. In particular, the present invention is intended to obtain a high-strength steel sheet having a high Yield Ratio (YR) exceeding 0.70, a TS (tensile strength) of 980MPa or more, excellent ductility, and excellent hole expansibility and bendability.
Background
In recent years, improvement of fuel efficiency of automobiles has become an important issue from the viewpoint of global environmental conservation. Therefore, there is an active trend toward weight reduction of a vehicle body itself by increasing the strength of a vehicle body material to achieve a reduction in thickness, but since increasing the strength of a steel sheet leads to a reduction in formability, development of a material having both high strength and high formability is desired. When a steel sheet having a TS of 980MPa or more is used as an automobile frame member, a high YR (yield ratio) is required for protecting an occupant in addition to high formability.
As a high-strength and excellent-ductility steel sheet, a high-strength steel sheet in which transformation is induced by working of retained austenite has been proposed. Such a steel sheet has a structure having retained austenite, and is easily formed by the retained austenite at the time of forming the steel sheet, while having high strength because the retained austenite is martensitic after forming.
For example, patent document 1 proposes a high-strength steel sheet having a tensile strength of 1000MPa or more and a total Elongation (EL) of 30% or more, which utilizes work-induced transformation of retained austenite, and has very high ductility. Such a steel sheet is produced by austenitizing a steel sheet containing C, Si, and Mn as basic components, quenching the steel sheet in a bainite transformation temperature range, and holding the steel sheet isothermally, that is, by so-called austempering. The C is enriched into austenite by the isothermal quenching treatment to form retained austenite, and a large amount of C exceeding 0.3% is required to be added in order to obtain a large amount of retained austenite. However, when the C concentration in steel is increased, the spot weldability is lowered, and particularly when the C concentration exceeds 0.3%, the spot weldability is remarkably lowered, and it is difficult to put the steel sheet into practical use as an automobile steel sheet. In addition, in the above patent documents, since the ductility of the high-strength thin steel sheet is mainly improved, hole expansibility and bendability are not considered.
In patent document 2, high Mn steel is used, and heat treatment in a dual phase region of ferrite and austenite is performed, thereby obtaining a high strength-ductility balance. However, in patent document 2, improvement of ductility due to enrichment of Mn into non-transformed austenite is not studied, and there is room for improvement of workability.
In patent document 3, Mn is enriched in the non-transformed austenite by applying heat treatment in a dual-phase region of ferrite and austenite in the medium Mn steel, thereby forming stable retained austenite and improving the total elongation. However, since the heat treatment time is short and the diffusion rate of Mn is slow, it is presumed that Mn is insufficiently enriched to achieve a yield ratio, hole expansibility, and bendability in addition to the elongation.
In patent document 4, in the use of Mn steel, a hot-rolled sheet is subjected to a long-time heat treatment in a dual-phase region of ferrite and austenite, thereby forming retained austenite having a large aspect ratio, which promotes the enrichment of Mn into non-transformed austenite, and improving the uniform elongation and the hole expansibility. However, the above documents have studied improvement of ductility and hole expansibility of a high-strength steel sheet enriched only by Mn, and have not studied improvement of yield ratio and bendability by controlling distribution of C and Mn in the second phase composed of retained austenite and martensite.
Documents of the prior art
Patent document
Patent document 1: japanese laid-open patent publication No. 61-157625
Patent document 2: japanese laid-open patent publication No. 1-259120
Patent document 3: japanese patent laid-open publication No. 2003-138345
Patent document 4: japanese patent No. 6123966
Disclosure of Invention
Problems to be solved by the invention
The present invention has been made in view of the above-described various current situations, and an object thereof is to provide a high-strength steel sheet having a high Yield Ratio (YR) exceeding 0.70, a TS (tensile strength) of 980MPa or more, and excellent formability, and a method for manufacturing the same. The formability described herein means ductility, hole expansibility, and bendability.
Means for solving the problems
The present inventors have made extensive studies from the viewpoint of the composition of the steel sheet and the production method thereof in order to solve the above problems, and have found the following.
Namely, the following are found: it is important that Mn is contained in an amount of 2.50 to 8.00 mass%, the composition of other alloying elements such as Ti is appropriately adjusted, and Ac is added as necessary after hot rolling1The cold rolling is performed while keeping the temperature within the range of not more than 1800 seconds. Then, at Ac3Keeping the temperature of the transformation point above-50 deg.C for 20s to 1800s, and cooling toA cooling stop temperature below the martensite start temperature. In this way, in the subsequent reheating step, film-like austenite which becomes nuclei of fine retained austenite having a large aspect ratio is generated. After the cooling is stopped, reheating the mixture to Ac1At least transformation point of Ac1A reheating temperature in a temperature range of +150 ℃ or lower, and then cooling the steel sheet after maintaining the steel sheet at the reheating temperature for 20 to 1800 seconds. Further, a zinc plating treatment is performed as necessary, and an alloying treatment is further performed at 450 ℃ or higher and 600 ℃ or lower as necessary. Then, the temperature is cooled to 100 ℃ or lower, and the temperature is maintained in a temperature range of over 100 ℃ to 400 ℃ or lower for 10 seconds or longer, and then the mixture is cooled. It is known that the following high-strength steel sheet having excellent formability can be produced by this method: in the steel structure, ferrite is 30-80% in terms of area ratio, tempered martensite is 3.0-35% in terms of area ratio, retained austenite is 8% or more, a value obtained by dividing the area ratio of the retained austenite of 2.0-1 [ mu ] m in terms of short axis by the area ratio of all the retained austenite is 0.3 or more, a value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in ferrite is 1.5 or more, a value obtained by multiplying the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in ferrite by the average of the retained austenite is 3.0 or more, and a value obtained by dividing the average C amount (mass%) in the retained austenite by the average C amount (mass%) in ferrite is 3.0 or more, and a value obtained by dividing the average C amount (mass%) in the retained austenite by the average Mn amount (mass%) in the retained austenite is less than 0.05, and the amount of diffusible hydrogen in the steel is 0.3 mass ppm or less.
The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] A high-strength steel sheet comprising:
the composition comprises the following components: contains, in mass%, C: 0.030% to 0.250% of Si: 0.01% or more and 3.00% or less, Mn: 2.50% -8.00%, P: 0.001% or more and 0.100% or less, S: 0.0001% or more and 0.0200% or less, N: 0.0005% or more and 0.0100% or less, Al: 0.001% to 2.000%, with the balance being Fe and unavoidable impurities; and
the following steel structure: wherein ferrite is 30 to 80% in terms of area ratio, tempered martensite is 3.0 to 35% in terms of area ratio, retained austenite is 8% or more, the area ratio of retained austenite having an aspect ratio of 2.0 to 1 μm or less in terms of short axis is 0.3 or more in terms of the area ratio of all retained austenite, the value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite is 1.5 or more, the value obtained by multiplying the value obtained by dividing the average Mn amount (mass%) in the ferrite by the average Mn amount (mass%) in the retained austenite by the average aspect ratio of the retained austenite is 3.0 or more, the value obtained by dividing the average C amount (mass%) in the retained austenite by the average C amount (mass%) in the ferrite is 3.0 or more, the value obtained by dividing the average C amount (mass%) in the retained austenite by the average Mn amount (mass%) in the retained austenite is less than 0.05,
the amount of diffusible hydrogen in the steel is 0.3 ppm by mass or less.
[2] The high-strength steel sheet according to [1], wherein the above-mentioned composition further contains, in mass%, a metal selected from the group consisting of Ti: 0.200% or less, Nb: 0.200% or less, V: 0.500% or less, W: 0.500% or less, B: 0.0050% or less, Ni: 1.000% or less, Cr: 1.000% or less, Mo: 1.000% or less, Cu: 1.000% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ta: 0.100% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, REM: 0.0050% or less.
[3] The high-strength steel sheet according to [1] or [2], further comprising a zinc plating layer on the surface.
[4] The high-strength steel sheet according to [3], wherein the zinc-plated layer is an alloyed zinc-plated layer.
[5] A method for producing a high-strength steel sheet having a steel structure comprising: a ferrite content of 30 to 80% in terms of area percentage, a tempered martensite content of 3 to 35% in terms of area percentage, a retained austenite content of 8% or more, an area percentage of a retained austenite content of 2.0 or more and a short axis content of 1 μm or less, a value obtained by dividing the area percentage of the total retained austenite content by the area percentage, a value obtained by dividing an average Mn content (mass%) in the retained austenite by an average Mn content (mass%) in the ferrite by the area percentage, a value obtained by multiplying a value obtained by dividing an average Mn content (mass%) in the retained austenite by an average Mn content (mass%) in the ferrite by the average aspect ratio of the retained austenite, a value obtained by dividing an average C content (mass%) in the retained austenite by an average C content (mass%) in the ferrite by the average aspect ratio, a value obtained by dividing an average C content (mass%) in the retained austenite by the average Mn content (mass%) in the retained austenite by the residual austenite by the average Mn content (mass%) in the ferrite by the mass%) is less than 0.05, the amount of diffusible hydrogen in the steel is 0.3 mass ppm or less,
in the manufacturing method,
to have [1]]Or [2]]The steel slab having the above composition is hot-rolled, then coiled at 300 to 750 ℃, cold-rolled, and Ac3Keeping the temperature of transformation point-50 ℃ or higher for 20s to 1800s, cooling to a cooling stop temperature of martensite transformation start temperature or lower, and reheating to Ac1At least transformation point of Ac1After a reheating temperature within a temperature range of +150 ℃ or lower, the steel sheet is kept at the reheating temperature for 20 seconds to 1800 seconds, and then cooled to 100 ℃ or lower, and further kept at a temperature in a range of more than 100 ℃ to 400 ℃ for 10 seconds or higher, and then cooled.
[6]Such as [5]]The method for producing a high-strength steel sheet, wherein Ac is added after coiling1The temperature range below the phase transition point is kept over 1800 s.
[7]Such as [5]]Or [6 ]]The method for producing a high-strength steel sheet, wherein Ac is used as the component1At least transformation point of Ac1The steel sheet is cooled after being kept at a temperature of 20 seconds to 1800 seconds within a range of +150 ℃ inclusive, and then is subjected to a galvanization treatment, and then is cooled to 100 ℃ inclusive.
[8] The method for producing a high-strength steel sheet as recited in [7], wherein the alloying treatment of the zinc-plated layer is performed in a temperature range of 450 ℃ to 600 ℃ after the zinc-plating treatment.
Effects of the invention
According to the present invention, a high-strength steel sheet having a high Yield Ratio (YR) exceeding 0.70, a TS (tensile strength) of 980MPa or more, and formability, particularly excellent in not only ductility but also hole expansibility and bendability, can be obtained. By applying the high-strength steel sheet obtained by the production method of the present invention to, for example, an automobile structural member, fuel efficiency improvement by weight reduction of a vehicle body can be achieved, and the industrial utility value is extremely high.
Detailed Description
The present invention will be specifically described below. The "%" indicating the content of the component element means "% by mass" unless otherwise specified.
(1) The reason why the composition of the steel is limited to the above range in the present invention will be described.
C: 0.030% to 0.250%
C is an element necessary for securing a low-temperature phase transformation phase such as tempered martensite to increase the strength. C is an element effective for improving the stability of retained austenite and improving the ductility of steel. When the C content is less than 0.030%, it is difficult to secure a desired area ratio of tempered martensite, and a desired strength cannot be obtained. In addition, it is difficult to secure a sufficient area fraction of retained austenite, and good ductility cannot be obtained. On the other hand, if C is added in excess of 0.250%, the area ratio of hard tempered martensite becomes too large, and micropores at grain boundaries of the tempered martensite increase during a hole expansion test, and propagation of cracks progresses, thereby reducing hole expandability. Further, the weld zone and the heat affected zone are significantly hardened, and the mechanical properties of the weld zone are reduced, so that the spot weldability, the arc weldability, and the like are also deteriorated. From such a viewpoint, the C content is set to 0.030% or more and 0.250% or less. Preferably 0.080% or more and 0.200% or less.
Si: 0.01% to 3.00%
Si is effective for ensuring good ductility because it improves the work hardening ability of ferrite. When the amount of Si is less than 0.01%, the effect of addition is poor, and therefore, the lower limit is set to 0.01%. However, excessive addition of Si exceeding 3.00% causes not only reduction in ductility and reduction in bendability due to embrittlement of steel, but also deterioration in surface properties due to occurrence of red rust and the like. In addition, a reduction in the quality of the plating layer results. Therefore, Si is set to 0.01% or more and 3.00% or less. Preferably 0.20% or more and 2.00% or less. More preferably 0.20% or more and less than 0.70%.
Mn: 2.50% or more and 8.00% or less
Mn is an extremely important additive element in the present invention. Mn is an element that stabilizes retained austenite, is effective for ensuring good ductility, and increases the strength of steel by solid-solution strengthening. In addition, Mn is effective for securing stable retained austenite enriched with Mn and obtaining good ductility. Such an effect is exhibited by setting the Mn content of the steel to 2.50% or more. However, if the Mn content exceeds 8.00% and the Mn content is excessively added, the area ratio of hard tempered martensite may become too large, and micropores at grain boundaries of the tempered martensite may increase during the hole expansion test, and propagation of cracks may progress, resulting in a decrease in hole expandability. In addition, the chemical conversion treatability and the quality of the plating layer are deteriorated. From such a viewpoint, the Mn content is set to 2.50% or more and 8.00% or less. Preferably 3.10% or more and 6.00% or less. More preferably 3.20% or more and 4.20% or less.
P: 0.001% or more and 0.100% or less
P is an element which has a solid-solution strengthening effect and can be added according to a desired strength. P is an element that promotes ferrite transformation and is therefore effective for composite organization. In order to obtain such an effect, the P content needs to be 0.001% or more. On the other hand, if the P content exceeds 0.100%, weldability deteriorates, and if the galvanized layer is alloyed, the alloying rate decreases, and the quality of the galvanized layer deteriorates. Therefore, the P amount is set to 0.001% or more and 0.100% or less, preferably 0.005% or more and 0.050% or less.
S: 0.0001% or more and 0.0200% or less
S segregates to grain boundaries to embrittle the steel during hot working, and exists as sulfides to reduce local deformability. Therefore, the S amount needs to be set to 0.0200% or less, preferably 0.0100% or less, and more preferably 0.0050% or less. However, the S content needs to be set to 0.0001% or more due to the restrictions in production technology. Therefore, the S amount is set to 0.0001% or more and 0.0200% or less, preferably 0.0001% or more and 0.0100% or less, and more preferably 0.0001% or more and 0.0050% or less.
N: 0.0005% or more and 0.0100% or less
N is an element that deteriorates the aging resistance of steel. In particular, if the N content exceeds 0.0100%, deterioration in aging resistance becomes significant. The smaller the amount of N, the more preferable, but the amount of N needs to be set to 0.0005% or more due to the restrictions in production technology. Therefore, the N amount is set to 0.0005% or more and 0.0100%, preferably 0.0010% or more and 0.0070% or less.
Al: 0.001% or more and 2.000% or less
Al is an element effective for enlarging the two-phase region of ferrite and austenite and reducing the annealing temperature dependence of mechanical properties, i.e., material stability. When the content of Al is less than 0.001%, the effect of addition is poor, so the lower limit is set to 0.001%. Further, Al is also an element that functions as a deoxidizer and is effective for the cleanliness of steel, and is preferably added in the deoxidation step. However, the addition of a large amount exceeding 2.000% increases the risk of cracking of the steel sheet during continuous casting, and decreases the manufacturability. From such a viewpoint, the Al content is set to 0.001% or more and 2.000% or less. Preferably 0.200% or more and 1.200% or less.
In addition, in addition to the above components, the composition may further contain, in mass%, a component selected from the group consisting of Ti: 0.200% or less, Nb: 0.200% or less, V: 0.500% or less, W: 0.500% or less, B: 0.0050% or less, Ni: 1.000% or less, Cr: 1.000% or less, Mo: 1.000% or less, Cu: 1.000% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ta: 0.100% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, REM (abbreviation of Rare Earth Metal): 0.0050% or less.
Ti: less than 0.200%
Ti is effective for precipitation strengthening of steel, and by increasing the strength of ferrite, the difference in hardness with the hard second phase (tempered martensite or retained austenite) can be reduced, and further excellent hole expansibility can be ensured. When Ti is added, it is preferably added in an amount of 0.005% or more. More preferably 0.010% or more. However, if the ratio exceeds 0.200%, the area ratio of hard tempered martensite may become too large, and micropores at grain boundaries of the tempered martensite may increase during the hole expansion test, and propagation of cracks may progress, thereby deteriorating the hole expandability. Therefore, when Ti is added, the amount of Ti added is set to 0.200% or less. Preferably, the content is set to 0.100% or less.
Nb: 0.200% or less, V: 0.500% or less, W: less than 0.500%
Nb, V, and W are effective for precipitation strengthening of steel. In addition, as with the effect of Ti addition, the strength of ferrite is increased, whereby the difference in hardness with the hard second phase (tempered martensite or retained austenite) can be reduced, and further excellent hole expansibility can be ensured. When Nb, V and W are added, they are preferably added in an amount of 0.005% or more, respectively. More preferably 0.010% or more. However, if Nb exceeds 0.200% and V, W exceeds 0.500%, the area fraction of hard tempered martensite becomes too large, and micropores at grain boundaries of the tempered martensite increase during a hole expansion test, and propagation of cracks progresses, and hole expandability decreases in some cases. Therefore, when Nb is added, the amount of Nb added is set to 0.200% or less, preferably 0.100% or less. When V, W is added, the amount of addition is set to 0.500% or less, preferably 0.300% or less.
B: 0.0050% or less
B has an action of suppressing the generation and growth of ferrite from austenite grain boundaries, and increases the strength of ferrite, thereby making it possible to reduce the difference in hardness with the hard second phase (tempered martensite or retained austenite), and to ensure further excellent hole expansibility. When B is added, it is preferably 0.0003% or more. More preferably 0.0005% or more. However, if it exceeds 0.0050%, moldability may be deteriorated. Therefore, when B is added, the amount of B added is set to 0.0050% or less. Preferably, the content is set to 0.0030% or less.
Ni: 1.000% or less
Ni is an element that stabilizes retained austenite and is effective for ensuring better ductility, and Ni is an element that increases the strength of steel by solid-solution strengthening. When Ni is added, it is preferably 0.005% or more. On the other hand, if the amount exceeds 1.000%, the area fraction of hard tempered martensite may become too large, and micropores at grain boundaries of the tempered martensite may increase during a hole expansion test, and crack propagation may progress, thereby deteriorating hole expandability. Therefore, when Ni is added, the amount of addition is set to 1.000% or less.
Cr: 1.000% or less, Mo: 1.000% or less
Cr and Mo have an effect of improving the balance between strength and ductility, and therefore may be added as needed. When Cr and Mo are added, they are preferably 0.005% or more, respectively. However, exceeding Cr: 1.000%, Mo: if the amount of the martensite is 1.000% and the martensite is excessively added, the area ratio of the hard tempered martensite may become excessively large, and micropores at grain boundaries of the tempered martensite may increase during the hole expansion test, and propagation of cracks may progress, thereby deteriorating the hole expandability. Therefore, in the case of adding these elements, the amounts thereof are set to 1.000% or less, respectively.
Cu: 1.000% or less
Cu is an element effective for strengthening steel, and may be added as needed. When Cu is added, it is preferably 0.005% or more. On the other hand, if the content exceeds 1.000%, the area ratio of the hard tempered martensite may become too large, and micropores at grain boundaries of the tempered martensite may increase during the hole expansion test, and propagation of cracks may progress, thereby deteriorating the hole expandability. Therefore, the amount of Cu added is set to 1.000% or less.
Sn: 0.200% or less, Sb: less than 0.200%
Sn and Sb are added as necessary from the viewpoint of suppressing decarburization of the surface layer of the steel sheet in a region of about several tens μm due to nitriding or oxidation of the surface of the steel sheet. Sn and Sb are effective for suppressing such nitriding and oxidation, preventing a reduction in the area ratio of tempered martensite on the steel sheet surface, and ensuring strength and material stability. When Sn and Sb are added, they are preferably 0.002% or more, respectively. On the other hand, if any of these elements is added in excess of 0.200%, the toughness may be lowered. Therefore, when Sn and Sb are added, the contents thereof are set to 0.200% or less, respectively.
Ta: less than 0.100%
Like Ti and Nb, Ta forms alloy carbide and alloy carbonitride, contributing to high strength. In addition, the following effects are considered to be obtained: since a part of the composite precipitates (Nb, Ta) (C, N) are formed as solid solutions in Nb carbide and Nb carbonitride, coarsening of the precipitates is significantly suppressed, and contribution to strength by precipitation strengthening is stabilized. Therefore, Ta is preferably contained. In the case where Ta is added, it is preferably 0.001% or more. On the other hand, even if Ta is excessively added, the precipitate stabilizing effect is saturated and the alloy cost is increased in some cases. Therefore, in the case of adding Ta, the content thereof is set to 0.100% or less.
Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, REM: 0.0050% or less
Ca. Mg, Zr, and REM are effective elements for spheroidizing the shape of the sulfide and improving the adverse effect of the sulfide on the hole expansibility. When these elements are added, they are preferably 0.0005% or more, respectively. However, excessive addition of more than 0.0050% may cause an increase in inclusions and the like, and surface and internal defects. Therefore, when Ca, Mg, Zr, and REM are added, the amounts of addition are set to 0.0050% or less, respectively.
The balance other than the above components being Fe and inevitable impurities.
(2) Next, the steel structure will be described.
Area ratio of ferrite: more than 30% and less than 80%
In order to ensure sufficient ductility, the area ratio of ferrite needs to be 30% or more. In order to ensure a tensile strength of 980MPa or more, the area ratio of the soft ferrite needs to be 80% or less. The ferrite referred to herein is polygonal ferrite, granular ferrite, or acicular ferrite, and is relatively soft and ductile ferrite. Preferably 40% or more and 75% or less.
Area ratio of tempered martensite: 3.0% to 35%
In order to ensure high local elongation, good hole expansibility and bendability, and a high yield ratio, tempered martensite needs to be 3.0% or more. In order to achieve high local elongation and good hole expansibility, bendability, and high yield ratio, the area ratio of tempered martensite needs to be set to 3.0% or more. In order to achieve a TS of 980MPa or more, the area ratio of tempered martensite needs to be 35% or less. Preferably 5.0% or more and 20% or less.
The area ratios of ferrite and tempered martensite can be determined as follows: after polishing a sheet thickness section (L section) parallel to the rolling direction of the steel sheet, etching was performed with a 3 vol% nital solution, 10 fields of view were observed at a magnification of 2000 times with an SEM (scanning electron microscope) at a position 1/4 in the sheet thickness (corresponding to a position 1/4 of the sheet thickness from the surface of the steel sheet in the depth direction), and using the obtained structure Image, the area ratios of the structures (ferrite and tempered martensite) in 10 fields of view were calculated with Image-Pro of Media Cybernetics, and the area ratios of the ferrite and tempered martensite were determined by averaging these values. In the structure image, ferrite shows a gray structure (matrix structure), and tempered martensite shows a gray internal structure inside white (quenched) martensite.
Area ratio of retained austenite: more than 8 percent
In order to ensure sufficient ductility, the area ratio of retained austenite needs to be 8% or more. Preferably 12% or more and 25% or less.
The area ratio of retained austenite is determined as follows: after polishing the steel sheet to a surface 0.1mm away from the 1/4 th position of the sheet thickness, the steel sheet was further polished by chemical polishing to 0.1mm, and the diffraction peaks of {200}, {220}, {311} plane of fcc iron and {200}, {211}, and {220} plane of bcc iron were measured for the surfaces obtained by the above-mentioned chemical polishing using CoK.alpha.rays using an X-ray diffractometer, and the 9 obtained integrated intensity ratios were averaged to determine the area ratio of retained austenite.
A value obtained by dividing the area ratio of retained austenite having an aspect ratio of 2.0 or more and a minor axis of 1 [ mu ] m or less by the area ratio of all the retained austenite is 0.3 or more
A value obtained by dividing the area ratio of retained austenite having an aspect ratio of 2.0 or more and a minor axis of 1 μm or less by the area ratio of all retained austenite of 0.3 or more is an important technical feature in the present invention. The retained austenite having an aspect ratio of 2.0 or more and a minor axis of 1 μm or less suppresses generation of voids during punching before the hole expanding step, and thus contributes to improvement of hole expandability. In order to ensure good hole expandability, it is necessary to increase the area ratio of the retained austenite having an aspect ratio of 2.0 or more and a minor axis of 1 μm or less while ensuring an area ratio sufficient to obtain high ductility of the retained austenite. Preferably 0.5 or more. The upper limit of the aspect ratio is preferably 15.0 or less. The lower limit of the minor axis is preferably 0.05 μm or more, which is the detection limit of EBSD.
Tempered martensite and retained austenite are identified by Phase Map of EBSD (Electron Backscattered Diffraction). The aspect ratio of retained austenite is calculated as follows: the length-to-diameter ratio of the retained austenite was calculated by plotting an ellipse circumscribing the retained austenite grains using Photoshop elements 13 and dividing the major axis length by the minor axis length.
Value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite: 1.5 or more
A value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite of 1.5 or more is an extremely important technical feature in the present invention. In order to ensure good ductility, a high area fraction of Mn-enriched stable retained austenite is required. Preferably 2.0 or more. The upper limit is not particularly limited, and is preferably 10.0 or less because the ductility is improved as the average Mn content in the retained austenite is higher.
The value obtained by multiplying the average aspect ratio of retained austenite by the value obtained by dividing the average Mn content (% by mass) in retained austenite by the average Mn content (% by mass) in ferrite is 3.0 or more
A value obtained by multiplying a value obtained by dividing the average Mn content (mass%) in the retained austenite by the average Mn content (mass%) in the ferrite by the average aspect ratio of the retained austenite by 3.0 or more is an extremely important technical feature. In order to ensure good ductility, it is necessary to have a large aspect ratio and a high area ratio of Mn-enriched stable retained austenite. Further, when a value obtained by multiplying a value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite by the average aspect ratio of the retained austenite is less than 3.0, generation of voids may become remarkable at the time of punching before the hole expanding step, and the hole expandability may be lowered. Preferably 4.0 or more. The upper limit is preferably 20.0 or less.
The value obtained by dividing the amount of C in retained austenite (mass%) by the average amount of C in ferrite (mass%): 3.0 or more
A value obtained by dividing the C content (mass%) in the retained austenite by the average C content (mass%) in the ferrite of 3.0 or more is an extremely important technical feature in the present invention. In order to ensure good ductility and bendability, it is necessary to increase the area ratio of stable retained austenite in which C is enriched. Preferably 5.0 or more. The upper limit is preferably 10.0 or less.
The value obtained by dividing the average C content (mass%) in the retained austenite by the average Mn content (mass%) in the retained austenite is less than 0.05
A value obtained by dividing the amount (mass%) of C in the retained austenite by the amount (mass%) of Mn in the retained austenite by less than 0.05 is an extremely important technical feature in the present invention. In order to ensure a high YR, Mn is more abundant than C in the retained austenite in which C and Mn are stable, and thus the stability of the retained austenite is increased and the yield stress becomes high. Preferably 0.02 or more and 0.04 or less.
The amounts of C and Mn in the retained austenite and ferrite were measured by cutting out a sample from a position 1/4 in the plate thickness using a three-Dimensional Atom Probe (3 DAP: 3Dimensional Atom Probe). First, a portion containing residual austenite and ferrite is cut out and processed into a needle-shaped sample by using a focused ion beam. The amount of Mn can be determined in atomic% by applying a voltage to the acicular sample with 3DAP, analyzing the C and Mn ions released at that time, and dividing the number of C and Mn atoms measured for each of the retained austenite and ferrite by the number of other total atoms. The above-described operation was performed on 30 random retained austenite grains and 30 ferrite grains in the measurement field, and the average value of the quantitative analysis results of the C, Mn amount was obtained. The amounts (mass%) of C and Mn in the retained austenite and ferrite were obtained by converting the C, Mn amount (atomic%) obtained above into mass%.
The steel structure of the present invention does not impair the effects of the present invention even when carbides such as quenched martensite, bainite, pearlite, and cementite are contained in an area ratio of 10% or less in addition to ferrite, tempered martensite, and retained austenite.
The amount of diffusible hydrogen in steel is 0.3 ppm by mass or less
The amount of diffusible hydrogen in steel of 0.3 mass ppm or less is an important technical feature in the present invention. In order to ensure high local elongation and good hole expansibility, the amount of diffusible hydrogen in the steel needs to be set to 0.3 mass ppm or less. The amount of diffusible hydrogen in steel is preferably in the range of 0.2 mass ppm or less. A test piece having a length of 30mm and a width of 5mm was cut out from the annealed sheet, and after removing the plating layer by grinding, the amount of diffusible hydrogen and the emission peak of diffusible hydrogen in the steel were measured. The peak was measured by Thermal analysis (TDS), and the temperature rise rate was set at 200 ℃ per hour. Hydrogen detected at 300 ℃ or lower is referred to as diffusible hydrogen.
The surface of the steel sheet may have a zinc plating layer. The zinc-plated layer may be an alloyed zinc-plated layer obtained by alloying.
(3) Next, the production conditions will be described.
Heating temperature of steel billet
Although not particularly limited, when the billet is heated, the heating temperature of the billet is preferably set to 1100 ℃ or higher and 1300 ℃ or lower. Since precipitates existing in the heating stage of the billet are present as coarse precipitates in the finally obtained steel sheet and do not contribute to strength, the heating temperature of the billet is preferably set to 1100 ℃ or higher from the viewpoint of redissolving Ti and Nb-based precipitates precipitated at the time of casting, removing bubbles, segregation, and the like in the surface layer of the billet, further reducing cracks and irregularities in the surface of the steel sheet, and achieving a smoother surface of the steel sheet. On the other hand, the heating temperature of the billet is preferably set to 1300 ℃ or lower from the viewpoint of reducing the scale loss accompanying the increase in the oxidation amount. More preferably, the temperature is set to 1150 ℃ or higher and 1250 ℃ or lower.
The billet is preferably produced by a continuous casting method in order to prevent macro-segregation, but may be produced by an ingot casting method, a thin slab casting method, or the like. In addition to the conventional method of once cooling to room temperature and then reheating after producing a billet, an energy saving process such as direct feed rolling or direct rolling in which the billet is charged into a heating furnace in a warm state without being cooled to room temperature or immediately rolled after being slightly held at temperature can be applied without any problem. Further, although the slab is roughly rolled into a thin slab under normal conditions, when the heating temperature is focused, it is preferable to heat the thin slab using a bar heater or the like before finish rolling from the viewpoint of preventing troubles during hot rolling.
Finish rolling exit side temperature of hot rolling: 750 ℃ or higher and 1000 ℃ or lower
The heated slab is hot-rolled by rough rolling and finish rolling to produce a hot-rolled steel sheet. At this time, when the finish rolling temperature exceeds 1000 ℃, the amount of oxide (scale) formed increases rapidly, the interface between the steel substrate and the oxide becomes rough, and the surface quality after pickling and cold rolling tends to deteriorate. Further, if hot-rolled scale remains locally after pickling, ductility and hole expansibility may be adversely affected. Further, the crystal grain size may become excessively large, resulting in a rough surface of the molded product during processing. On the other hand, when the finish rolling temperature is less than 750 ℃, the rolling load may be increased, the reduction ratio of austenite in a non-recrystallized state may be increased, and an abnormal texture may be developed. As a result, in-plane anisotropy of the final product becomes remarkable, uniformity of material (material stability) is impaired, and not only does this deteriorate ductility itself in some cases. In addition, the aspect ratio of the retained austenite is decreased, and ductility and hole expansibility are sometimes decreased. Therefore, the temperature of the finish rolling outlet side of the hot rolling is preferably set to 750 ℃ to 1000 ℃. Preferably, the temperature is set to 800 ℃ or higher and 950 ℃ or lower.
Coiling temperature after hot rolling: 300 ℃ or higher and 750 ℃ or lower
When the coiling temperature after hot rolling exceeds 750 ℃, the crystal grain size of ferrite in the hot rolled sheet structure increases, the aspect ratio of the retained austenite of the final annealed sheet decreases, and the hole expansibility decreases. On the other hand, when the coiling temperature after hot rolling is lower than 300 ℃, the hot-rolled sheet strength increases, the rolling load during cold rolling increases, or a defect in sheet shape occurs, and therefore productivity decreases. Therefore, it is necessary to set the coiling temperature after hot rolling to 300 ℃ to 750 ℃. Preferably, the temperature is set to 400 ℃ or higher and 650 ℃ or lower.
In addition, the rough rolled plates may be joined to each other during hot rolling to continuously perform finish rolling. Further, the rough rolled sheet may be temporarily wound. In addition, in order to reduce the rolling load during hot rolling, part or all of the finish rolling may be performed by lubrication rolling. From the viewpoint of uniformizing the shape and uniformizing the material quality of the steel sheet, it is also preferable to perform lubrication rolling. When the lubrication rolling is performed, the friction coefficient during the lubrication rolling is preferably set to 0.10 or more and 0.25 or less.
The hot rolled steel sheet thus manufactured is optionally subjected to pickling. Pickling is preferably performed because it can remove oxides on the surface of the steel sheet and further improve chemical conversion treatability and coating quality. Further, the pickling may be performed once or in a plurality of times.
At Ac1The temperature below the phase change point is kept to be more than 1800s
At Ac1The steel sheet to be used for the subsequent cold rolling is softened while being kept in a temperature range of not more than 1800 seconds at the transformation point, and therefore, the cold rolling is preferably performed. In excess of Ac1When the steel sheet is kept in the temperature range of the transformation point, Mn may be concentrated in austenite, and after cooling, hard quenched martensite and retained austenite are formed, and the steel sheet may not be softened. In addition, in excess of Ac1When the alloy is held in the temperature range of the transformation point, the alloy may form massive quenched martensite and retained austenite, and may continue to have a massive structure in the subsequent heat treatment, thereby decreasing the aspect ratio and reducing the hole expansibility. In addition, even in Ac1When the temperature is kept in the range of 1800 seconds or less within the temperature range of the transformation point or less, it may be difficult to remove the strain after hot rolling, and the steel sheet may not be softened.
The heat treatment method may be either continuous annealing or batch annealing. Further, after the heat treatment, the steel sheet is cooled to room temperature, but the cooling method and the cooling rate are not particularly limited, and any of furnace cooling in batch annealing, air cooling, and gas jet cooling, spray cooling, and water cooling in continuous annealing may be used. In addition, in the case of performing the acid washing treatment, a conventional method can be employed.
Cold rolling
The resulting steel sheet was subjected to cold rolling. The cold rolling reduction is not limited, but is preferably 15% or more and 80% or less. By performing cold rolling in this range, a desired structure sufficiently recrystallized can be obtained, and the properties are improved.
At Ac3The temperature range of transformation point-50 ℃ is kept between 20s and 1800s
Below Ac3When the alloy is maintained at a temperature of-50 ℃ C, Mn is concentrated in austenite, and martensite transformation does not occur during cooling, and a core of retained austenite having a large aspect ratio cannot be obtained. As a result, in the subsequent annealing step, retained austenite is formed from grain boundaries, and the retained austenite having a small aspect ratio increases, so that a desired structure cannot be obtained. Even at Ac3The temperature range of-50 ℃ or higher is maintained for less than 20s, and the process is not carried outSufficient recrystallization does not lead to a desired structure, and ductility is reduced. In addition, Mn surface enrichment for ensuring the quality of the subsequent plating layer is not sufficiently performed. On the other hand, if the steel sheet is kept at a temperature of more than 1800 seconds, Mn surface enrichment may saturate, and hard tempered martensite and retained austenite may increase on the surface of the steel sheet after the final annealing treatment, thereby deteriorating bendability.
Cooling to a cooling stop temperature below the martensite start temperature
When the cooling stop temperature exceeds the martensite transformation start temperature, the amount of quenched martensite at which transformation occurs is small, and a core of retained austenite having a large aspect ratio cannot be obtained. As a result, in the subsequent annealing step, retained austenite is formed from grain boundaries, and the retained austenite having a small aspect ratio increases, so that a desired structure cannot be obtained. Further, the Mn enrichment in the retained austenite is reduced, and thus a high YR may not be obtained in some cases. Preferably, the martensite transformation starting temperature is-250 ℃ or higher and the martensite transformation starting temperature is-50 ℃ or lower.
Is heated again to Ac1At least transformation point of Ac1After a reheating temperature within a temperature range of +150 ℃ or lower, the resultant composition is maintained at the reheating temperature for 20 seconds to 1800 seconds
At Ac1At least transformation point Ac1The temperature range of the phase transition point +150 ℃ or lower is maintained for 20 seconds or more and 1800 seconds or less, which is an extremely important technical feature of the invention in the present invention. Below Ac1When the alloy is held in the temperature range of the transformation point and in the condition of less than 20 seconds, the carbide formed by the temperature increase is not completely melted, and it is difficult to ensure a sufficient area ratio of tempered martensite and retained austenite after the final annealing treatment, and the strength is lowered. In addition, C and Mn are insufficiently enriched into the retained austenite, and ductility and bendability are reduced. In excess of Ac1In the temperature range of +150 ℃ as the transformation point, the area ratio of martensite increases, and Mn is enriched and saturated in austenite, so that sufficient area ratio of retained austenite sites is not obtained, and ductility decreases. In addition, the retained austenite forms in a lump, and the aspect ratio and hole expansibility are reduced. Preferably Ac1Transformation point +100 ℃ or lower. Furthermore, inWhen the amount exceeds 1800 seconds, austenite growth in the short axis direction is promoted, and the aspect ratio is decreased, thereby reducing hole expandability. Further, since C is enriched into the retained austenite, it is difficult to obtain a value obtained by dividing the average C amount (mass%) in the desired retained austenite by the average Mn amount (mass%) in the retained austenite, and it is difficult to secure a high YR. When the zinc plating treatment is subsequently performed, cooling is performed once. The cooling stop temperature before galvanization is preferably 350 ℃ to 550 ℃.
Galvanizing treatment
When the hot dip galvanizing treatment is performed, the annealing treatment (reheating to Ac) is performed1At least transformation point of Ac1A reheating temperature within a range of +150 ℃ or lower, and then holding at the reheating temperature for 20 seconds or more and 1800 seconds or less) followed by immersing the steel sheet, which has been cooled to a temperature of at least the temperature of the galvanizing bath by gas jet cooling, furnace cooling, or the like, in the galvanizing bath at 440 ℃ or more and 500 ℃ or less, performing a hot galvanizing treatment, and then adjusting the amount of deposit of the plating layer by gas wiping or the like. In the hot dip galvanizing, a galvanizing bath having an Al content of 0.08% to 0.30% is preferably used. In addition to the hot dip galvanizing treatment, a method such as an electrogalvanizing treatment may be used.
When the alloying treatment of the galvanized layer is performed, the alloying treatment of the galvanized layer is performed in a temperature range of 450 ℃ to 600 ℃ after the galvanizing treatment. When the alloying treatment is performed at a temperature exceeding 600 ℃, the non-transformed austenite is transformed into pearlite, and a desired area ratio of retained austenite cannot be secured, and ductility may be reduced. Therefore, when the alloying treatment of the galvanized layer is performed, the alloying treatment of the galvanized layer is preferably performed in a temperature range of 450 ℃ to 600 ℃.
Cooling to below 100 deg.C
By cooling to 100 ℃ or lower, a sufficient amount of tempered martensite capable of securing tensile strength can be formed after the final annealing treatment. Further, it is preferable to cool the glass to about 20 ℃ to about 50 ℃ before the subsequent step of holding the glass in a temperature range of more than 100 ℃ to 400 ℃ for 10 seconds or more. In the case where the above-mentioned galvanizing treatment or alloying treatment of the galvanized layer is performed, the treatment is performed before the step of cooling to 100 ℃.
Keeping the temperature of the mixture for more than 10s in a temperature range of more than 100 ℃ and less than 400 DEG C
As the final heat treatment (final annealing treatment), it is important in the present invention to maintain the temperature in the range of more than 100 ℃ and 400 ℃ or less for 10 seconds or more. When the steel is held in a temperature range of 100 ℃ or less or for less than 10 seconds, tempered martensite having a sufficient area fraction cannot be obtained, and diffusible hydrogen in the steel is not released from the steel sheet, so that not only hole expansibility is lowered but also bendability is reduced. On the other hand, when the steel is held at a temperature in the range exceeding 400 ℃, the retained austenite is decomposed, so that a sufficient area fraction of the retained austenite is not obtained, and the ductility of the steel is lowered.
The conditions of the other production method are not particularly limited, and the annealing is preferably performed by a continuous annealing facility from the viewpoint of productivity. In addition, a series of processes such as annealing, hot Galvanizing, and alloying of a galvanized layer is preferably performed by a hot Galvanizing Line CGL (Continuous Galvanizing Line).
The "high-strength steel sheet" and the "high-strength galvanized steel sheet" having a galvanized layer on the surface of the high-strength steel sheet may be skin-rolled for the purpose of shape correction, surface roughness adjustment, and the like. The reduction ratio of skin pass rolling is preferably in the range of 0.1% to 2.0%. If the content is less than 0.1%, the effect is small and the control is difficult, so that the lower limit of the preferable range is set. In addition, when it exceeds 2.0%, productivity is remarkably lowered, so that it is set as the upper limit of a preferable range. It should be noted that the skin pass rolling may be performed on-line or off-line. Further, the skin pass rolling at the target reduction ratio may be performed at one time, or may be performed in a plurality of times. Various coating treatments such as resin coating and grease coating may be performed.
Examples
Will have the formula shown in Table 1The steel composed of the components and the balance of Fe and unavoidable impurities is melted in a converter and made into a billet by a continuous casting method. The obtained slabs were reheated to 1250 ℃, and then high-strength cold-rolled steel sheets (CR) having a thickness of 1.0mm to 1.8mm were obtained under the conditions shown in table 2. Further, the hot dip galvanized steel sheet (GI) was obtained by galvanizing, and the hot dip galvanized steel sheet was alloyed to obtain an alloyed hot dip galvanized steel sheet (GA). As for the hot dip galvanizing bath, a zinc bath containing 0.19 mass% of Al was used for the hot dip galvanized steel sheet (GI), and a zinc bath containing 0.14 mass% of Al was used for the galvannealed steel sheet (GA), and the bath temperature was set at 465 ℃. The amount of deposit was set to 45g/m per surface2(double-sided plating), the Fe concentration in the plating layer is adjusted to 9 mass% or more and 12 mass% or less for GA. The cross-sectional steel structure, tensile properties, hole expandability, and bendability of the obtained steel sheets were examined, and the results are shown in tables 3, 4, and 5.
Figure BDA0003605784270000221
Ms point and Ac point of martensite transformation initiation temperature1Transformation point and Ac3The transformation point was determined using the following equation.
Martensite transformation Start temperature Ms point (. degree. C) — 550-350 × (% C) -40 × (% Mn) -10 × (% Cu) -17 × (% Ni) -20 × (% Cr) -10 × (% Mo) -35 × (% V) -5 × (% W) +30 × (% Al)
Ac1Phase Change Point (. degree.C.) 751-16 × (% C) +11 × (% Si) -28 × (% Mn) -5.5 × (% Cu) -16 × (% Ni) +13 × (% Cr) +3.4 × (% Mo)
Ac3Phase Change Point (. degree.C.) 910-
Wherein, (% C), (% Si), (% Mn), (% Ni), (% Cu), (% Cr), (% Mo), (% V), (% Ti), (% V), (% W), and(% Al) represent the content (mass%) of each element.
[ Table 2]
Figure BDA0003605784270000241
Underlined section: are shown outside the scope of the invention.
CR: cold-rolled steel sheet (no plating), GI: hot-dip galvanized steel sheet (galvannealed steel sheet), GA: alloyed hot-dip galvanized steel sheet
[ Table 3]
Figure BDA0003605784270000251
Underlined section: are shown outside the scope of the invention.
F: ferrite, TM: tempered martensite, RA: retained austenite
The tensile test was carried out in accordance with JIS z 2241 (2011) using a JIS5 test piece obtained by cutting a sample so that the tensile direction was perpendicular to the rolling direction of the steel sheet, and TS (tensile strength), EL (total elongation), YS (yield stress), or YR (yield ratio) was measured. Here, YR is a value obtained by dividing YS by TS. In the present invention, it is judged that mechanical properties are good when YR exceeds 0.70 and the mechanical properties satisfy the following.
TS: EL is more than or equal to 20 percent under the condition of more than 980MPa and less than 1080MPa
TS: EL is more than or equal to 16 percent under the condition of more than 1080MPa and less than 1180MPa
Hole expandability was performed in accordance with JIS Z2256 (2010). Each of the obtained steel sheets was cut into 100mm × 100mm, a hole having a diameter of 10mm was punched out with a gap of 12% ± 1%, a punch of a 60 ° cone was pressed into the hole with a die having an inner diameter of 75mm pressed with a pressing force of 9 tons, the hole diameter at the time of crack occurrence limit was measured, a limit hole expansion ratio λ (%) was determined according to the following formula, and the hole expansibility was evaluated based on the value of the limit hole expansion ratio.
Limiting hole expansion ratio λ (%) { (D)f-D0)/D0}×100
Wherein D isfThe pore diameter (mm) at the time of crack generation, D0Is the initial holeDiameter (mm). In the present invention, the following case is determined to be preferable for each TS range.
TS: under the condition of 980MPa above and 1080MPa below, the lambda is more than or equal to 15 percent
TS: the lambda is more than or equal to 12 percent under the condition that the pressure is more than 1080MPa and less than 1180MPa
In the Bending test, a Bending test piece having a width of 30mm and a length of 100mm was cut out from each annealed steel sheet so that the rolling direction was the Bending axis (Bending direction), and the Bending test piece was measured by the V-block method according to JIS Z2248 (1996). A test was performed at each bending radius at a pressing speed of 100 mm/sec so that the presence or absence of a crack was determined by a stereomicroscope on the outer side of the bending portion, and the minimum bending radius at which no crack occurred was defined as the limit bending radius R. In the present invention, it is determined that the bendability of the steel sheet is good when the minimum bend R/t of 90 DEG V-bend is not more than 2.5 (t: the sheet thickness of the steel sheet).
The high-strength steel sheets of the examples of the present invention all had a TS of 980MPa or more and had excellent formability. On the other hand, in the comparative example, at least one of the characteristics of YR, TS, EL, λ and bendability was poor.
[ Table 4]
Figure BDA0003605784270000271
Underlined section: are shown outside the scope of the invention.
F: ferrite, M: quenched martensite, RA: retained austenite, BF: bainitic ferrite,
P: pearlite, θ: carbide (cementite, etc.)
[ Table 5]
Figure BDA0003605784270000281
Underlined section: are shown outside the scope of the invention.
Industrial applicability
According to the present invention, a high-strength steel sheet having a Yield Ratio (YR) exceeding 0.70 and having a TS (tensile strength) of 980MPa or more and excellent formability can be obtained. By applying the high-strength steel sheet of the present invention to, for example, an automobile structural member, fuel efficiency improvement due to weight reduction of a vehicle body can be achieved, and the industrial utility value is very high.

Claims (8)

1. A high-strength steel sheet comprising:
the composition comprises the following components: contains, in mass%, C: 0.030% to 0.250% of Si: 0.01% or more and 3.00% or less, Mn: 2.50% or more and 8.00% or less, P: 0.001% or more and 0.100% or less, S: 0.0001% or more and 0.0200% or less, N: 0.0005% or more and 0.0100% or less, Al: 0.001% to 2.000%, with the balance being Fe and unavoidable impurities; and
the following steel structure: wherein ferrite is 30 to 80% in terms of area ratio, tempered martensite is 3.0 to 35% in terms of area ratio, retained austenite is 8% or more, the area ratio of retained austenite having an aspect ratio of 2.0 to 1 μm or less in terms of short axis is 0.3 or more in terms of the area ratio of all retained austenite, the value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite is 1.5 or more, the value obtained by multiplying the value obtained by dividing the average Mn amount (mass%) in the ferrite by the average Mn amount (mass%) in the retained austenite by the average aspect ratio of the retained austenite is 3.0 or more, the value obtained by dividing the average C amount (mass%) in the retained austenite by the average C amount (mass%) in the ferrite is 3.0 or more, the value obtained by dividing the average C amount (mass%) in the retained austenite by the average Mn amount (mass%) in the retained austenite is less than 0.05,
the amount of diffusible hydrogen in the steel is 0.3 mass ppm or less.
2. The high-strength steel sheet according to claim 1, wherein the composition further contains, in mass%, a component selected from the group consisting of Ti: 0.200% or less, Nb: 0.200% or less, V: 0.500% or less, W: 0.500% or less, B: 0.0050% or less, Ni: 1.000% or less, Cr: 1.000% or less, Mo: 1.000% or less, Cu: 1.000% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ta: 0.100% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, REM: 0.0050% or less.
3. The high-strength steel sheet according to claim 1 or 2, further comprising a zinc plating layer on the surface.
4. The high strength steel sheet according to claim 3, wherein the galvanized layer is an alloyed galvanized layer.
5. A method for producing a high-strength steel sheet having a steel structure comprising: a ferrite content of 30 to 80% in terms of area percentage, a tempered martensite content of 3 to 35% in terms of area percentage, a retained austenite content of 8% or more, an area percentage of a retained austenite content of 2.0 or more and a short axis content of 1 μm or less, a value obtained by dividing the area percentage of the total retained austenite content by the area percentage, a value obtained by dividing an average Mn content (mass%) in the retained austenite by an average Mn content (mass%) in the ferrite by the area percentage, a value obtained by multiplying a value obtained by dividing an average Mn content (mass%) in the retained austenite by an average Mn content (mass%) in the ferrite by the average aspect ratio of the retained austenite, a value obtained by dividing an average C content (mass%) in the retained austenite by an average C content (mass%) in the ferrite by the average aspect ratio, a value obtained by dividing an average C content (mass%) in the retained austenite by the average Mn content (mass%) in the retained austenite by the residual austenite by the average Mn content (mass%) in the ferrite by the mass%) is less than 0.05,
the amount of diffusible hydrogen in the steel is 0.3 mass ppm or less,
in the manufacturing method,
a steel slab having the composition of claim 1 or 2, hot-rolled, coiled at 300 ℃ to 750 ℃, cold-rolled, and Ac3Keeping the temperature of transformation point-50 ℃ or above for 20s to 1800s, and cooling to martensiteCooling to a temperature below the transformation starting temperature, and reheating to Ac1At least transformation point of Ac1After a reheating temperature within a temperature range of +150 ℃ or lower, the steel sheet is held at the reheating temperature for 20 seconds to 1800 seconds, and then cooled to 100 ℃ or lower, and further held at a temperature within a range of more than 100 ℃ to 400 ℃ for 10 seconds or higher, and then cooled.
6. The method for producing a high-strength steel sheet according to claim 5, wherein Ac is added after coiling1The temperature range below the phase transition point is kept over 1800 s.
7. The method for producing a high-strength steel sheet according to claim 5 or 6, wherein Ac is added to the steel sheet1At least transformation point of Ac1The steel sheet is cooled after being kept at a temperature of 20 seconds to 1800 seconds within a range of +150 ℃ inclusive, then subjected to a galvanization treatment, and then cooled to 100 ℃ inclusive.
8. The method for producing a high-strength steel sheet according to claim 7, wherein the alloying treatment of the zinc-plated layer is performed at a temperature range of 450 ℃ to 600 ℃ after the zinc-plating treatment.
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EP3447160A1 (en) * 2016-04-19 2019-02-27 JFE Steel Corporation Steel plate, plated steel plate, and production method therefor
JP6338038B1 (en) * 2017-11-15 2018-06-06 新日鐵住金株式会社 High strength cold-rolled steel sheet
WO2019111029A1 (en) * 2017-12-05 2019-06-13 Arcelormittal Cold rolled and annealed steel sheet and method of manufacturing the same
WO2019188642A1 (en) * 2018-03-30 2019-10-03 Ｊｆｅスチール株式会社 High-strength steel sheet and method for manufacturing same

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TWI831367B (en) * 2022-09-08 2024-02-01 大田精密工業股份有限公司 High strength nickel chromium molybdenum steel and manufacturing method the same
CN115491608A (en) * 2022-09-29 2022-12-20 桂林理工大学 One kind (Cr, fe) 7 C 3 TiC composite reinforced medium manganese steel and preparation method thereof
CN115491608B (en) * 2022-09-29 2023-08-15 桂林理工大学 A kind of (Cr, fe) 7 C 3 TiC composite reinforced medium manganese steel and preparation method thereof

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US20220396847A1 (en) 2022-12-15
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